Wet chemical etchants

ABSTRACT

Silicon-germanium-based compositions comprising silicon, germanium, and carbon (i.e., Si--Ge--C), methods for growing Si--Ge--C epitaxial layer(s) on a substrate, etchants especially suitable for Si--Ge--C etch-stops, and novel methods of use for Si--Ge--C compositions are provided. In particular, the invention relates to Si--Ge--C compositions, especially for use as etch-stops and related processes and etchants useful for microelectronic and nanotechnology fabrication.

The invention was made with U.S. Government support under (i) Phase ISBIR N00014-93-C-0114 awarded by the Office of Naval Research (BMDO),(ii) F49620-93-C-0018 awarded by AFOSR (DARPA), and (iii) DMR-9115680awarded by the National Science Foundation, and the Government hascertain rights in the invention.

This application is a continuation of copending application Ser. No.08/336,949, filed Nov. 10, 1994.

BACKGROUND

The present invention relates to silicon-germanium-based compositionscomprising silicon, germanium and carbon (Si--Ge--C), methods forgrowing Si--Ge--C epitaxial layer(s) on a substrate, etchants especiallysuitable for Si--Ge--C etch-stops, and novel methods of use forSi--Ge--C compositions. In particular, the present invention relates toSi--Ge--C compositions, especially for use as etch-stops, relatedprocesses and etchants useful for microelectronic and nanotechnologyfabrication. The present application is a continuation of, claimspriority to and incorporates by reference the entire disclosure of U.S.application Ser. No. 08/336,949, filed on Nov. 10, 1994.

In etching we remove a film or layer from a substrate, in some instancesdefining the layer to be removed by photolithography. One way to etch isto immerse the substrate in a bath of some chemical that attacks thefilm. Preferably, the chemical should react with and etch the film orlayer in a smooth and reproducible manner, producing soluble productsthat can be carried away from the substrate. In particular, an idealetchant will not attack any layer underneath the film being etched, sothat the etch process will be self-limiting. Unfortunately, the etchingis often not self-limiting and therefore goes below the desired depth.Etch-stops are designed to address this problem.

Semiconductors have the interesting property that when they are alloyedwith certain elements, the rate of wet chemical etch in the alloy willvary from that of the unalloyed semiconductor. Alloys with differentetch rates can be used to cause etching to slow at a pre-definedinterface. Typically, a layer with a particular composition etches at aknown rate in an etchant A second adjacent layer may etch at a differentrate because it has a different composition. The layer with the loweretch rate is often referred to as the etch-stop layer.

Etch-stops are used to fabricate devices for a wide variety ofapplications. Membranes and diaphragms formed via etch-stops are used insensors such as pressure transducers, as elements in experimental x-raylithography systems, as windows for high energy radiation, and as lowthermal mass supports for microcalorimeter and bolometric radiationdetectors. Additional uses for selective etch-stops are inmicromachining applications such as accelerometers, gears, micro-beams,miniature fluid lines, pumps and valves, and in flow sensors. Anotherapplication with a potentially large commercial market is in fabricatingsilicon-on-insulator (SOI) substrates by the bond-and-etch-backsilicon-on-insulator (BESOI) process.

Silicon-based selective chemical etch-stop layers such as Si--B, Si--Ge,Si--Ge--B, Si--P, and Si--As have major problems and disadvantages whichare overcome by the present invention. The disadvantages can beillustrated by examining examples pertaining to the commonly usedSi--Ge--B etch-stops. First, the specially doped layer (e.g., Si--Ge--B)and the lightly doped silicon have limited selectivity. Selectivity isdefined as the etch rate of lightly doped silicon divided by the etchrate of the etch-stop layer, or in some cases its reciprocal asdiscussed below. Limited selectivity increases the manufacturing cost bycreating a need for tightly controlled, and sometimes labor-intensiveprocessing to prevent the etch from going beyond the intended depth.This problem is exacerbated when fabricating the thin layers that arerequired for submicron electronic devices.

Certain chemical solutions etch a lightly doped silicon layer morerapidly than a heavily doped layer. For this purpose lightly doped meansless than approximately 1E17 dopant atoms per cm³, and heavily dopedmeans more than approximately 1E19 dopant atoms per cm³. For example, 21weight percent (wt %) potassium hydroxide in H₂ O (KOH-H₂ O) at about70° C. etches the (100) plane of lightly doped silicon rapidly(approximately 1 micrometer per minute), but the etch rate becomes slow(less than 0.01 micrometer per minute), making possible selectiveetching, as the boron concentration in the silicon increases to morethan about 5E19 atoms per cm³.

The conventional formulations of these etchants have serious problemswhen used with the above-mentioned etch-stop layers. For example, KOH-H₂O, an inexpensive etchant, is prone to producing a rough surface when itetches silicon. Ethylenediamine pyrocatechol in water (EDP-H₂ O)provides somewhat better etch selectivity and is less prone todeveloping surface roughness than KOH-H₂ O, but has limited applicationin that it emits extremely toxic vapors and is relatively expensive.Another etchant, cesium hydroxide in water (CsOH-H₂ O), can providesmoother surfaces than KOH-H₂ O for conventional etch-stop layers but iseven more expensive than EDP-H₂ O.

Surface roughness arises from the anisotropic etch properties of thesolutions that preferentially etch lightly doped silicon. Thesesolutions etch certain crystallographic directions in the materialfaster than other directions. For example, a chemical solutionconsisting of 21 wt % KOH-H₂ O at 70° C. will rapidly etch the (100)plane of lightly doped silicon, but only slowly etch the (111) plane.This leads to the etched surface being rough as illustrated in FIGS.1A-B. As shown in FIG. 1A, a solution of KOH-H₂ O (11) will rapidly etchlightly doped silicon (12). If a small particle such as particle (13)adheres to the surface of the lightly doped silicon (12), the etch ratewill be locally retarded under the particle (13). Slow etching planes(14) on the (111) plane will form as the particle (13) is undercut bythe etch solution. This leads to the formation of a slow-etching pyramidunder the particle (13). If the etch selectivity is not sufficientlyhigh, these pyramids will propagate into the etch-stop layer (15), shownin FIG. 1A as peaks (16), resulting in a rough surface.

Another problem with conventional etch-stop compositions, especiallythose containing a high boron concentration, is they leave an insolublestaining residue on the surface of the substrate. This residue bothroughens and contaminates the substrate surface. Increasing the KOHconcentration of the etch solution, for example, from 21 wt % to 40 wt %will eliminate the surface staining, but will also substantiallydecrease the etch selectivity.

In contrast to the above etchants, conventional formulations of 1:3:8and 1:3:12 parts by volume of HF-HNO₃ -CH₃ COOH (HNA) etch lightly dopedsilicon somewhat less rapidly than heavily doped silicon and aretherefore used preferentially to remove etch-stop layers. In this case,selectivity is defined as the etch rate of the etch-stop layer dividedby the etch rate of the lightly doped silicon. The above formulations ofHNA have major drawbacks, including relatively low selectivity, andselectivity decreasing rapidly with time during etching due to reductionof the HNO₃ to HNO₂.

Still another problem with conventional etch-stop compositions is thatthe impurity which provides the etch selectivity is also a donor oracceptor dopant in the silicon. Thus, for example, when a Si--Ge--Betch-stop is used to fabricate a BESOI substrate, boron diffusing outfrom the etch-stop layer during a bonding anneal causes unwantedelectrically active dopant in the device layer. This problem isillustrated in FIGS. 2A-B which show concentration profiles of boron andgermanium as a function of depth in a substrate layer, an etch-stoplayer, and a device layer, before a bonding anneal (FIG. 2A) and afterthe anneal (FIG. 2B). FIG. 2A illustrates the boron (21) and germanium(22) concentration profiles in the substrate (23), the etch-stop layer(24), and in the device layer (25) after epitaxial layer growth andbefore the bonding anneal. FIG. 2B shows the changed boron (26) andgermanium (27) concentration profiles after the bonding anneal. Becauseboron diffuses through the material faster than germanium during thebonding anneal, its profile is broadened such that significant"diffusion tails" extend from the etch-stop layer into the substrate(23) and the device layer (25). In a BESOI structure, the borondiffusion tail causes an unacceptable level of electrically activedopant to exist in the device layer (25).

There are reports in the literature of Si--Ge--C layer fabrication.However, to the best of applicants' knowledge, none of the existingprocesses for forming Si--Ge--C are suitable for producing Si--Ge--Clayers as part of a large scale manufacturing process. Feijoo et al.,Etch Stop Barriers in Silicon Produced by Ion Implantation ofElectrically Non-Active Species, Journal of the Electrochemical Society(1992) describe silicon layers implanted with silicon, germanium, andcarbon at doses between 1E14 and 3E16 ions/cm² and energies between 35and 200 keV and testing them as etch-stop barriers in an EDP-H₂ O basedsolution (p. 2309, Abstract). When ions are implanted in this range ofdose and energies, the lattice structure is damaged. Feijoo states theresults obtained indicate that the effectiveness of the etch-stop isinfluenced (i.e., improved) by both the implantation damage and thechemical interaction between the implanted ions and the defectivecrystal (Abstract). The resulting damage greatly restricts the number ofuseful commercial applications for Feijoo's etch-stop barriers.Accordingly, Feijoo's methods and results are substantially differentfrom the present invention.

U.S. Pat. No. 4,885,614 to Furukawa et al., Semiconductor Device withCrystalline Silicon-Germanium-Carbon Alloy, describes another process ofproducing a silicon-germanium-carbon alloy film principally by molecularbeam epitaxy, but also by plasma enhanced chemical vapor deposition(CVD), photoenhanced CVD, microwave-excited CVD, thermal CVD andmetal-organic CVD methods. Molecular beam epitaxy (MBE) might providegood crystalline quality, but it is a slow and expensive process. Withregard to a description of the thermal CVD process (col. 10, lines37-43) Furukawa describes that the surface of a silicon substrate wascleaned and the temperature thereof adjusted to 650° C. Gaseous SiH₄,GeH₄ and CH₄ were allegedly introduced into a reactor so as to give atotal pressure of 100 torr. Thus, a Si--Ge--C alloy film was purportedlyformed on the substrate by a thermal CVD method. However, applicantsbelieve methane at the stated process temperature is far too stable tofunction as a carbon source for thermal CVD formation of thesilicon-germanium-carbon film. Further, Furukawa et al. do not recognizeor discuss the use of Si--Ge--C as an etch-stop.

Regolini et al., Growth and characterization of strain compensatedSi_(1-x-y) Ge_(x) C_(y) epitaxial layers, Materials Letters (1993)describe metal-organic chemical vapor deposition (MOCVD) for fabricatingepitaxial Si--Ge--C layers with less than 1 atomic percent carbon, whichis far less than desirable for etch-stops. Further, there is no mentionin the Regolini et al. publication of using Si--Ge--C as an etch-stop.

SUMMARY OF INVENTION

This invention describes a silicon-based etch-stop which is composed ofan alloy containing silicon, germanium and carbon (Si--Ge--C). In oneembodiment, this alloy demonstrates etch selectivity of 8000 to 1compared with lightly doped silicon. High etch selectivity isdemonstrated for both the case in which the Si--Ge--C epitaxial layeracts as the etch-stop, and for the opposite case in which the Si--Ge--Clayer is preferentially removed.

This invention further describes a method by which the etch-stop layeris grown epitaxially and with excellent crystal quality by CVD onto asilicon substrate. One or more high quality, undoped or deliberatelydoped epitaxial silicon-based layers can be grown before and/or afterthe etch-stop layer.

This invention further describes an innovative range of etch solutioncompositions which, combined with the unique composition of theetch-stop layer, enhance the selectivity for removal of the etch-stoplayer by more than two orders of magnitude over previous processes foretch-stop layer removal, and compared with previous formulations,provide an etch rate which varies far less with time. Further advantageswill be apparent from the following description of the preferredembodiments.

BRIEF DESCRIPTION OF THE DRAWINGS

FIGS. 1A-B are cross-sections illustrating surface roughness ofconventional etch-stop layers.

FIGS. 2A-B are graphs illustrating the dopant diffusion tail that iscommon with conventional etch-stop layers.

FIG. 3 is a perspective view illustrating a SOI substrate.

FIGS. 4A-F are cross-sections illustrating a BESOI process using anetch-stop layer.

FIGS. 5A-B are cross-sections illustrating reduced surface roughnesswith a Si--Ge--C etch-stop layer.

FIGS. 6A-B are graphs illustrating the absence of an electrically activedopant tail with a Si--Ge--C etch-stop layer.

FIGS. 7A-F are ball-and-stick representations of crystal structureillustrating the effect of carbon and germanium atoms on stress insilicon alloys.

FIG. 8 is a graph illustrating the superior etch selectivity ofSi--Ge--C in KOH-H₂ O compared with Si--Ge--B.

FIG. 9 is a graph illustrating the superior etch selectivity ofSi--Ge--C in HNA compared with Si--Ge--B.

FIGS. 10A-C are cross-sections illustrating a process for making amembrane using an etch-stop layer.

FIGS. 11A-C are cross-sections illustrating a process for making acantilever beam using an etch-stop layer.

FIG. 11D is a perspective view illustrating the cantilever beam formedusing the process of FIGS. 11A-C.

FIGS. 12A-B illustrate 2.0 MeV He²⁺ RBS random and (100) channeledspectra for: (a) sample A-6; (b) sample B-6. The dechanneling observedclose to the Si--Ge--C/Si interface suggested the presence of misfitdislocations.

FIG. 13 illustrates 2.0 MeV He²⁺ RBS random and (100) channeled spectrafor: (a) sample C-6. The Si--Ge--C was not even grown epitaxially on the(100)Si substrate.

FIG. 14 illustrates 2.0 MeV He²⁺ RBS random and (100) channeled spectrafor sample D-6. The Si--Ge--C was initially grown epitaxially. Asubsequent c-a (crystalline-to-amorphous) phase transformation occurred.

FIG. 15 illustrates cross-sectional TEM micrograph and correspondingcalibrated carbon SIMS profile of sample D-6. The Si--Ge--C can beroughly divided into three regions: crystalline, highly defective, andamorphous region.

DETAILED DESCRIPTION OF THE PREFERRED EMBODIMENTS

Example: BESOI Fabrication:

One important application of the present invention is in fabricatingBESOI substrates. BESOI is used for advanced integrated circuitapplications in which the device semiconductor layer (31) shown in FIG.3 is isolated from the base wafer (32) by an insulating layer (33). Thefollowing description illustrates advantages of the invention.

An embodiment of a BESOI process is illustrated in FIGS. 4A-F. As shownin FIG. 4A, BESOI fabrication starts with two substrates, a base wafer(32) of a silicon wafer which can be either a lightly or heavily dopedsilicon wafer doped with either P-type or N-type dopant, such as boronor phosphorous, and a device wafer (41) of lightly doped silicon waferof either of the same P-type or N-type dopants, and preferably a P-typedopant, such as boron. A Si--Ge--C layer (42) is grown epitaxially ontothe device wafer (41), and a device epitaxial silicon layer (31) isgrown epitaxially onto the Si--Ge--C layer. The Si--Ge--C layer (42) is50-500 nm thick, is preferably 75-200 nm thick, and is most preferablyabout 100 nm thick. The Si--Ge--C layer (42) includes 2-6 atomic percentcarbon, preferably 4-5 atomic percent carbon, and most preferably about4.5 atomic percent carbon, and includes 18-65 atomic percent germanium,preferably 35-50 atomic percent germanium, and most preferably about 40atomic percent germanium. The device silicon epitaxial layer (31) is500-2000 nm thick, is preferably 100-150 nm thick, and is mostpreferably about 120 nm thick, and contains electrically active dopantof the P-type or N-type in the range of 1E13-1E19 atoms per cm³,preferably 1E14-1E18 atoms per cm³, and most preferably about 1E15 atomsper cm³.

In the next step, thermal oxide layers (43, 44) shown in FIG. 4B aregrown by a conventional technique such as exposing the surface of eitherone or preferably both of the wafers (32, 41) to water vapor at800-1000° C. for about 5-300 minutes. W. S. Ruska, MicroelectronicProcessing (1987) describes such conventional techniques and is herebyincorporated by reference in its entirety. The base wafer oxide layer(43) thickness is 50-1000 nm, preferably 100-500 nm, and most preferablyabout 200-300 nm. The device wafer oxide (44) thickness is 20-100 nm,preferably 40-70 nm, and most preferably about 50-60 nm. The oxide layerthickness is preferably less on the device wafer (41) than on the basewafer (32) so that the thermal oxidation step will not diffuse orprecipitate the carbon out of the Si--Ge--C layer.

The next step is to invert the device wafer (41) and to bond it to thebase wafer (32) by contacting the respective oxide surfaces (44, 43)firmly together to form an oxide layer (33) as shown in FIG. 4C.

FIG. 4D illustrates that in the next step most of the device wafer (41)is removed by lapping or grinding, and optionally polished by aconventional technique to form a sacrificial silicon layer (47). Thethinning is preferably stopped when approximately 10 to 50 micrometersof the sacrificial layer (47) remains over the etch-stop layer (42).

FIG. 4E illustrates the structure after the first selective etch step inwhich the final 10 to 50 micrometers of the sacrificial layer (47)(shown in FIG. 4D) is removed using a chemical solution such as 10 to 45wt % and preferably 21 wt % KOH-H₂ O, and in a temperature range of from50 to 100° C., and preferably at about 70° C., which will etch thesacrificial layer (47) (shown in FIG. 4D), but will substantially stopetching when it reaches the etch-stop layer (42).

FIG. 4F illustrates the result of the final etch step in which theetch-stop layer (42) shown in FIG. 4E is removed using a chemicalsolution such as HNA at room temperature which etches the Si--Ge--Clayer rapidly, but essentially stops etching when it reaches the lightlydoped epitaxial device layer (31). The final structure, shown also inFIG. 3, is the lightly doped epitaxial device layer (31) separated fromthe substrate (32) by the insulating oxide (33).

FIGS. 5A-B illustrate how the present invention solves the surfaceroughness problem involved in BESOI processing by the conventionalmethod shown in FIG. 1B. FIG. 5A again shows a solution of KOH-H₂ O (11)will rapidly etch lightly doped silicon (12). Again, if a small particle(13) adheres to the surface of the lightly doped silicon (12), the etchrate is locally retarded under the particle (13), and the slow etchingplanes (14) on the (111) plane form as the particle (13) is undercut bythe etch solution. However, in the present invention, the resultingpyramidal defect will no longer propagate into the etch-stop layer (15)because of the improved selectivity of the etch-stop layer. As shown inFIG. 5B, the surface roughness (51) is substantially reduced when theetch front reaches the etch-stop layer (15) due to the pyramidal defectsbeing substantially etched away before the etch front significantlyproceeds into the etch-stop layer (15).

Chemical vapor deposition of an epitaxial Si--Ge--C layer on silicon:

One aspect of the present invention is that we have discovered a methodto grow Si--Ge--C layers epitaxially by a commercially viable CVD methodonto silicon at temperatures that are low enough to "quench" carbon intothe epitaxial layer. Low temperature is important because whereas thegermanium-silicon system exhibits complete solid solubility, carbon isnearly insoluble in silicon. The maximum solubility of carbon in siliconis about 1 part per million at the melting point of silicon. Thus, toincorporate several atomic percent carbon into silicon the epitaxiallayer growth takes place at a low enough temperature so that the carbonatoms are essentially immobile in the silicon lattice. In practice, thegrowth temperature is preferably less than about 800° C. Numerouspublications on the CVD growth of silicon carbide suggest that for mosthydrocarbon gases a growth temperature greater than 1100° C. is requiredto decompose the hydrocarbon precursor gases to supply a gas source ofcarbon. In the present invention, we discovered that cyclopropane orethylene can provide a source of carbon that is sufficiently reactive todeposit carbon-containing layers in a CVD process at temperatures as lowas 600° C.

We have further discovered that certain novel processes significantlyimprove the ability to grow high quality Si--Ge--C layers. For example,the following process produces Si--Ge--C epitaxial layers with 42 atomicpercent germanium, 5 atomic percent carbon, and a low defect density. Alightly doped, P-type (100) 200 mm diameter prime silicon substrate iscleaned and loaded into a CVD reactor using a robotic transfer arm. Asuitable reactor for this application is an Epsilon One, Model E-2,manufactured by ASM America, Phoenix, Ariz. Initially, the reactortemperature is held in the range from 25-900° C., preferably 800-900°C., and most preferably about 850° C. Hydrogen gas is introduced intothe reactor at preferably atmospheric pressure, or at reduced pressureof, for example, 5-100 torr, at a flow rate of 10-100 standard litersper minute (slm), preferably 15-40 slm, most preferably about 20 slm.Next, the substrate is heated to 1000-1200° C., preferably to 1100-1190°C., and most preferably to about 1150° C. The reactor is most preferablyheld at about 1150° C. for 30 to 600 seconds, preferably 60 to 300seconds, and most preferably about 90 seconds to remove the native oxidelayer from the substrate. The substrate is then cooled to 900° C. in theH₂ atmosphere before the deposition of silicon. At 900° C., 5 to 200standard cubic centimeters per minute (sccm), preferably 10 to 50 sccm,and most preferably about 20 sccm of SiH₂ Cl₂ (dichlorosilane) isintroduced into the reactor atmosphere to initiate deposition of theundoped or lightly doped epitaxial silicon. Deposition continues as thesubstrate is cooled further, and at 700° C. a low flowrate (e.g., 0.1sccm) of GeH₄ (germane) is added to the deposition atmosphere. Thesubstrate is cooled still further, and at 550-750° C., preferably600-700° C., and most preferably about 625° C. the low flowrate of GeH₄is increased over a span of 15 seconds to a final flowrate of from 1.0to 20.0 sccm, preferably 1.4 to 5.0 sccm, and most preferably about 1.8sccm. Just as the GeH₄ flowrate reaches its final value, C₃ H₆(cyclopropane) is added at a low flowrate (e.g. 25 sccm) to thedeposition atmosphere. The low flowrate of cyclopropane is increasedover a span of 15 seconds to a final flowrate of from 50 to 200 sccm,preferably 75 to 150 sccm, and most preferably about 100 sccm. All finalgas flowrates are then held constant for 1 to 100 minutes, preferably 5to 10 minutes, and most preferably about 6.5 minutes while a Si--Ge--Clayer of approximately 200 nm thickness is grown. Following depositionthe reactor is purged with 20 slm of pure H₂ for about 45 seconds.Afterward the robotic transfer arm removes the substrate from thereactor. In an alternative embodiment, an alkene such as ethylene can besubstituted for cyclopropane in a similar process.

The reasons for the above process improvements are as follows. Thesurface of silicon forms a tenacious silicon dioxide (SiO₂) when exposedto air. This oxide is removed from silicon during heating in hydrogen toa high temperature such as 1150° C., but the oxide tends to reform whenthe silicon surface is exposed to trace amounts of oxygen or water vaporat a lower temperature. By starting the CVD silicon growth at 900° C.,and by adding a small concentration of GeH₄ to the gas at 700° C., weprevent trace amounts of oxygen and water vapor in the depositionatmosphere from causing oxide nuclei to form as the temperature islowered to the level that is favorable for Si--Ge--C deposition. As thisprotective Si--Ge layer is deposited the GeH₄ concentration is held lowand ramped up quickly to keep the thickness and stress of the layer wellbelow the critical thickness for formation of misfit dislocations. TheC₃ H₆ flowrate is ramped from a low value to the final value for atotally different reason. In developing this process we discovered thatthe best crystalline quality is obtained by starting the Si--Ge--Cdeposition with a low concentration of hydrocarbon and then ramping theflowrate up until the desired carbon concentration is obtained. Thereason may be related to the low solubility of carbon in silicon.Starting the Si--Ge--C deposition with a low surface concentration ofcarbon may be important in establishing a steady-state growth process inwhich carbon is incorporated in non-equilibrium concentrations into thesilicon-germanium lattice.

An economical epitaxial layer growth process is important if theepitaxial single Si--Ge--C layers are to be used for commercialapplications. As will be demonstrated in the examples which follow, theprocess which we discovered can be operated in a commercial CVDepitaxial reactor at rates sufficiently high that the Si--Ge--C layerscan be produced at prices comparable to present epitaxial siliconlayers. A suitable reactor for this application is again the EpsilonOne, Model E-2, manufactured by ASM America, Phoenix, Ariz. The reactorcan be configured for single wafer automated processing usingvacuum-compatible load lock ports and a nitrogen-purged robotic wafertransfer chamber. The load lock ports and the wafer transfer chambereffectively isolate the process chamber from atmospheric contamination.Evacuation of the load lock ports is important to prevent air fromentering the process chamber and to remove adsorbed moisture from thesurfaces of the substrates before processing.

High purity materials are important for low-temperature, defect-freeproduction of epitaxial Si--Ge--C layers. Hydrogen and nitrogen liquidsources having impurity levels less than 10 parts per billion (ppb) canbe obtained from Air Products & Chemicals, Inc., in Allentown, Pa.Dichlorosilane gas of Ultraplus grade can be obtained from Praxair, inKingman, Ariz.; and a 1:99 volume ratio and a 10:90 volume ratio ofgermane:hydrogen of Megabit grade can be obtained from SolkatronicChemicals, in Fairfield, N.J. The 10:90 volume ratio ofcyclopropane:helium of semiconductor grade can be obtained from AirProducts & Chemicals. The hydrogen, nitrogen, germane and cyclopropanegases can be further purified by point-of-use Nanochem purifiers made bySemi-Gas Systems, in Santa Clara, Calif.

Additional processes and embodiments for the formation of thinheteroepitaxial films of Si--Ge--C grown on silicon substrates using CVDand C₂ H₄ gas as a carbon source at the relatively low temperature of625° C. are described as Chemical Vapor Deposition Of HeteroepitaxialSi_(1-x-y) Ge_(x) C_(y) Films On (100)Si Substrates.

Thin heteroepitaxial films of Si_(1-x-y) Ge_(x) C_(y) have been grown on(100)Si substrates using atmospheric pressure Chemical Vapor Deposition(CVD) at 625° C. The crystallinity, composition and microstructure ofthe Si--Ge--C films were characterized using Rutherford backscatteringspectrometry, secondary-ion-mass spectrometry and cross-sectionaltransmission electron microscopy. The crystallinity of the films wasvery sensitive to the flow rate of C₂ H₄ which served as the C source.Films with up to 2% C were epitaxial with good crystallinity and veryfew interfacial defects. Between 800 and 900 sccm of 10% C₂ H₄ in He,the C content increased dramatically from 2% to 10% and the as-grownfilms changed from crystalline to amorphous. In order to establishdeposition conditions for the crystalline-amorphous phasetransformation, one Si--Ge--C film was deposited as the 10% C₂ H₄ flowwas increased linearly from 500 to 1500 sccm during growth. When the Ccontent reached ˜4%, the film developed considerable stacking defectsand disorder, and at around 11% C, the film became amorphous.

The significant successes of semiconductor electronic devices areclosely connected to the concept of band gap engineering. The band gapof Si-based materials has been successfully engineered by alloying withGe and by forming strained layer superlattices. See R. People, J. C.Bean, D. V. Lang, A. M. Sergent, H. L. Stromer, K. W. Wecht, R. T.Lynch, and K. Baldwin, Appl. Phys. Lett. 45, 1231 (1984), which isincorporated by reference in its entirety. The band gap decreasesmonotonically as the Ge concentration increases, although compressivestrain in pseudomorphic Si_(1-x) Ge_(x) grown epitaxially on (100)Sisubstrates causes stability problems and limits the film thickness. Inmany applications it is desirable to have a wider band gap than that ofpure Si. Carbon in its diamond form is an elemental group-IV insulator(wide band gap of 5.5 eV) with a lattice parameter of 0.3545 nm, muchsmaller than that of Si (0.5431 nm) or Ge (0.5646 nm). Addingsubstitutional carbon into Si--Ge alloy layers may increase the band gap(see Richard A. Soref, J. Appl. Phys. 70, 2470 (1991), which isincorporated by reference in its entirety) and, at the same time, resultin strain compensation. See S. Furukawa, H. Etoh, A. Ishizaka, and T.Shimada, U.S. Pat. No. 4,885,614, (1989), which is incorporated byreference in its entirety. See K. Eberl, S. S. Iyer, and F. K. LeGoues,Appl. Phys. Lett. 64, 739 (1994), which is incorporated by reference inits entirety. In fact, strain in Si_(1-x-y) Ge_(x) C_(y) alloy layerscan be adjusted over a wide range from compressive all the way totensile, depending on the Ge/C concentration ratio. If Vegard's law isapplied to the Si--Ge--C ternary system, then we expect compensation of8.2 atomic % Ge by 1 atomic % C in the pseudomorphic Si_(1-x-y) Ge_(x)C_(y) film. See S. S. Iyer, K. Eberl, M. S. Goorsky, F. K. LeGoues, J.C. Tsang, Appl. Phys. Lett. 60, 356 (1992), which is incorporated byreference in its entirety. Recently, strain compensation by sequentialimplantation of C in Si--Ge alloy layers has been investigated byforming Si_(1-x-y) Ge_(x) C_(y) layers embedded in Si. See J. W. Strane,H. J. Stein, S. R. Doyle, S. T. Picraux, and J. W. Mayer. Appl. Phys.Lett. 63, 2786 (1993), which is incorporated by reference in itsentirety. See A. Fukami, K. Shoji, T. Nagano, and C. Y. Yang, Appl.Phys. Lett. 57, 2345 (1990), which is incorporated by reference in itsentirety. However, since the maximum solubility limit of C in Si is2×10⁻³ atomic %, precipitation of silicon carbide has been a severeproblem, particularly since higher temperatures are required to grow thestructure and repair implantation damage. This section is focused on thegrowth of pseudomorphic Si_(1-x-y) Ge_(x) C_(y) layers on (100)Sisubstrates using Chemical Vapor Deposition (CVD).

Si_(1-x-y) Ge_(x) C_(y) layers were epitaxially grown on (100)Si in anEPSILON ONE, model E-2 single wafer, automated CVD reactor using SiH₂Cl₂ as a Si source, GeH₄ as a Ge source, and C₂ H₄ as the C source atatmospheric pressure. Point of use purifiers on the gas lines providepurity levels in the parts-per-billion range in the process chamber.Nitrogen-purged load locks exclude air from the system. Four differentsamples were grown at 625° C. at various C₂ H₄ flows while the flows ofthe other two precursors were held constant. The growth conditions aresummarized in Table 1 shown below.

Rutherford backscattering spectrometry (RBS) in the channeling mode andthe ¹² C(α,α)¹ 2C elastic resonance reaction at 4.265 MeV were used todetermine the composition, thickness and crystalline quality of thegrown layer. The distribution of carbon within the films wascharacterized by secondary-ion-mass spectroscopy (SIMS), and themorphology of the grown layers was examined by cross-sectionaltransmission electron microscopy (TEM).

                  TABLE 1    ______________________________________    Sample growth conditions in CVD system.                                    10% C.sub.2 H.sub.4                                           10% C.sub.2 H.sub.4           Deposition        GeH.sub.4                                    in He  in He           Temperature                     SiH.sub.2 Cl.sub.2                             1% in H.sub.2                                    start  end    Sample (°C.)                     (sccm)  (sccm) (sccm) (sccm)    ______________________________________    A-6    625       20      180    504    504    B-6    625       20      180    800    800    C-6    625       20      180    900    900    D-6    625       20      180    500    1500    ______________________________________

FIGS. 12(a) and (b) show the 2.0 MeV He²⁺ RBS random and (100) channeledspectra for the as-grown A-6 and B-6 samples, respectively. The measuredvalues for χ_(min) ranged from 0.05 to 0.08 for sample A-6 and 0.13 to0.22 for sample B-6, with the higher values obtained close to theSi--Ge--C/Si interface in both samples. A significant yield enhancementwas observed in the Si substrate region of the spectrum. Thisenhancement was caused by dechanneling, which suggested the presence ofdefects at the Si--Ge--C/Si interface. Complementary TEM observationsindicate that the majority of these defects are simple edgedislocations. The carbon content of these samples was measured by the ¹²C(α,α)¹ 2C elastic resonance reaction using a 4.28 MeV beam of He⁺⁺ions. It was found that the fractions of Ge and C in the Si_(1-x-y)Ge_(x) C_(y) layer were x=0.37 and y=0.012 for sample A-6 and x=0.36 andy=0.02 for sample B-6. The SIMS carbon profiles for these samples showedthat the carbon was distributed uniformly in the alloy layer.

FIG. 13 shows the 2.0 MeV He²⁺ RBS random and (100) channeled spectrafor the as-grown sample C.6. The overlapping of the random and alignedspectra that exists in the Si--Ge--C region indicates that this layerwas not grown epitaxially on the (100)Si substrate. The fraction of C inthe alloy, as measured by resonance analysis, was y=0.10 while the Gefraction was x=0.27. SIMS analysis showed that carbon was distributeduniformly in the alloy layer, and cross-sectional TEM showed that theSi--Ge--C layer was amorphous. It was observed that for a relativelysmall increase in the C₂ H₄ flow (from 800 to 900 sccm) the carboncontent increased dramatically and the layer went from crystalline toamorphous. Thus, under these growth conditions, it appears that there isa critical level of C₂ H₄ flow that changes the growth from epitaxial toamorphous, while the incorporated carbon increases from 2 atomic % to 10atomic %. Determination of this critical flow level may be veryimportant, since this parameter affects directly the maximumconcentration of carbon that can be incorporated into the single crystalSi--Ge matrix during the epitaxial growth that takes place in this CVDsystem.

In order to investigate this critical flow level, another sample wasgrown (sample D-6) using the same conditions (shown in Table I), exceptthat the C₂ H₄ flow was increased linearly from 500 sccm at thebeginning of growth to 1500 sccm at the end. FIG. 14 shows the 2.0 MeVHe²⁺ RBS random and (100) channeled spectra for this sample. We observethat the channeled spectrum overlaps the random spectrum for both Si andGe signals at the region near the surface, indicating the presence of anamorphous layer, while for the region of the spectrum closer to theSi--Ge--C/Si interface, the yield of the channeled spectrum decreasessharply. These observations reveal that at the beginning of the growthprocess the Si--Ge--C was grown epitaxially, but that acrystalline-amorphous phase transformation occurred sometime during thegrowth process. The random spectrum shows that the Ge concentrationdecreases monotonically from the interface to the surface while the Siconcentration does not change significantly.

A cross-sectional TEM micrograph and calibrated carbon SIMS profile ofsample D-6, shown in FIG. 15, corroborate the RBS results. Based onthese results, the Si--Ge--C layer can be roughly divided into threeregions. Immediately adjacent to the Si--Ge--C/Si interface there is acrystalline region, ˜90 nm thick, with a roughly linear increase incarbon concentration up to ˜4 atomic %. The next region, 30-40 nm thick,is highly defective and contains many lattice defects such as stackingfaults and twins. The carbon concentration in this region increasessharply from 4 to ˜10 atomic %. Careful analysis of lattice-fringes inthis region using optical diffraction methods reveals no unexpectedspacings, indicating in particular that no silicon carbide has beenformed. The third region near the surface is completely amorphous, ˜80nm thick, and the carbon concentration is ˜10-12 atomic %.

The concentration of the alloy obtained at the crystalline/defectiveregion interface is Si₀.62 Ge₀.34 C₀.04 (shown in FIG. 4), which isprobably the highest C content that can be achieved without defects inthe crystalline layer under these CVD conditions. The Ge:C ratio at thispoint is 8.5:1, which is very close to the fully strain-compensatedvalue (8.2:1) that was calculated using Vegard's law. See S. S. Iyer, K.Eberl, M. S. Goorsky, F. K. LeGoues, J. C. Tsang, Appl. Phys. Lett. 60,356 (1992), which is incorporated by reference in its entirety. ThisGe:C ratio decreases dramatically in the defective region to a minimumvalue of 2.9:1 at the defective region/amorphous interface. Thereappears to be some correlation between the relative Ge:C composition andthe occurrence of lattice defects, suggesting that tensile strain couldplay a role in their production. However, this interpretation may be toosimple-minded, since the strain around the individual C atoms will bemuch greater than around Ge atoms.

In spite of the fact that the C₂ H₄ flow was increased linearly duringgrowth, the carbon profile, as depicted by SIMS analysis, has asigmoidal shape. This result indicates that carbon atoms can beincorporated into the Si--Ge matrix of the grown layer much more easilywhen it is in an amorphous state, i.e., the C₂ H₄ precursor is morereactive when it comes into contact with an amorphous substrate relativeto a crystalline one. The GeH₄ precursor behaves in the opposite manner,since the Ge concentration decreases during growth. The reactivity ofthe SiH₂ Cl₂ precursor, however, did not change.

In summary, CVD of Si_(1-x-y) Ge_(x) C_(y) alloy layers on (100)Sisubstrates is reported. Alloy layers with up to 2 atomic % of carbonhave been grown epitaxially, and have good crystallinity. Above acritical level of C₂ H₄ flow, when the Ge:C ratio of the deposited layeris higher than 8.5:1, the grown layer becomes increasingly defective,and when Ge:C<˜3:1, an amorphous layer is eventually formed.

This work was supported by AFOSR (DARPA) Award F49620-93-C-0018.Electron microscopy was conducted at the Center for High ResolutionElectron Microscopy supported by the National Science Foundation underGrant DMR-9115680.

Advantages of Si--Ge--C being isoelectronic:

An additional advantage of the Si--Ge--C etch-stop over the mostcommonly used Si--Ge--B etch-stop is that all three components of theSi--Ge--C alloy are isoelectronic, i.e., none supplies an electricallyactive dopant. Thus, if there is slight diffusion of one of the elementsinto the silicon device layer, the electrically active dopantconcentration will not be affected. Further, since germanium and carbonaffect the band structure and band gap of silicon, it will beadvantageous for some applications to leave the Si--Ge--C etch-stoplayer in place to become part of the device structure. FIGS. 6A-Billustrate the point. FIG. 6A illustrates carbon (61) and germanium (62)concentration profiles over a substrate (63), on which an epitaxialSi--Ge--C etch-stop layer (64) and a lightly doped device layer (65)have been epitaxially grown before a bonding anneal. This structurerepresents the starting point for a BESOI process, which can be comparedto the Si--Ge--B structure shown in FIGS. 2A-B. In FIG. 6A, the boronconcentration (66) is deliberately low throughout the concentrationprofile. The substrate (63) is shown with a boron concentration of about1E17 atoms per cm³, and the epitaxial layers (64, 65) are grown with aboron concentration of about 2E15 atoms per cm³. FIG. 6B shows theconcentration profiles after the bonding anneal. The carbon (67) andboron (68) profiles broaden during the anneal due to diffusion. Incontrast to the Si--Ge--B etch-stop in FIG. 2B, the carbon (69) in thedevice layer (65) of FIG. 6B does not contribute electrical carriers andthe substrate boron diffusion tail (70), residing in the etch-stop layer(64), does not contribute electrical carriers to the device layer (65).

Advantages of Si--Ge--C permitting strain control:

The present invention also permits strain control in the Si--Ge--Clayer. Germanium and carbon atoms both differ in size from silicon. Thefourfold covalent atomic radius of germanium is about 5% larger thansilicon, whereas carbon is about 50% smaller. The effects of germaniumand carbon addition to silicon are shown schematically in FIGS. 7A-F. Inundoped silicon, there is naturally-occurring spacing (71) between thesilicon atoms, represented as solid dots in FIG. 7A. When germaniumatoms are added to silicon, the lattice parameter or atomic spacing (72)in the crystalline structure is increased by the larger germanium atomsrepresented by the large white dots in FIG. 7B. By contrast, when carbonatoms represented by the small grey dots are added to silicon, theatomic spacing (73) is decreased as shown in FIG. 7C. As shown in FIG.7D, when a Si--Ge epitaxial layer represented by the upper three rows ofsolid and white dots is grown onto a silicon substrate represented bythe lower five rows of solid dots, a compressive in-plane stress arisesas the Si--Ge layer is forced to assume the atomic spacing (71) of theunstrained silicon substrate. The Si--Ge lattice expands in a directionnormal to the silicon substrate to form spacing (74) causing a tensilestrain in that direction. Conversely, when a Si--C layer is grownepitaxially onto a silicon substrate, the Si--C layer is under anin-plane tensile strain and a compressive strain in the direction normalto the silicon substrate. Thus, when either a Si--Ge or a Si--C layer isgrown epitaxially onto silicon, the growing layer experiences a stressthat increases with the thickness of the layer and with theconcentration of Ge or C. At some critical thickness the stress exceedsthe yield stress, the point at which dislocations nucleate at theepitaxial layer-substrate interface and propagate to form atwo-dimensional defect array (75) as illustrated in FIG. 7E for the caseof Si--Ge. This array of misfit dislocations allows the epitaxial layerto assume its unstrained atomic spacing (72), and in doing so relievesthe stress in the epitaxial layer. If, however, the epitaxial layerincludes Ge and C in a ratio defined by: (atomic radius Si-atomic radiusC)/(atomic radius Ge-atomic radius Si), about 9:1 Ge:C, the epitaxiallayer will have similar atomic spacing (71) to undoped silicon and befree of stress as shown in FIG. 7F. In this case, the epitaxiallayer-substrate interface remains free of misfit dislocations.

In different applications, it will be advantageous to grow Si--Ge--Clayers that (1) are strain-free, or (2) have strain that is grown-in toa pre-defined level, or (3) have lattice parameter misfit at so high alevel as to cause misfit dislocations to form. Fabrication of acantilever beam in an accelerometer application is an example in whichthe layer should be grown stress-free so that the beam will not bebowed. On the other hand, when a Si--Ge--C layer is to be made into amembrane, a slightly tensile stress, created by increasing thecarbon-to-germanium ratio, is desirable so that the membrane will besmooth and "tight." For semiconductor devices, both strain andcomposition are means by which the band gap and band structure can bevaried for particular applications. At the same time, it is usuallyimportant for electronic applications that the epitaxial layers, whetherstrained or unstrained, must be free of defects including misfitdislocations. For some mechanical applications, on the other hand, itmay be desirable to create an array of misfit dislocations. Suchdislocations can strengthen the silicon by a process known as "workhardening."

The above-referenced interface misfit stress can be measured in any ofseveral ways. One is to measure the interatomic spacing of a Si--Ge--Clayer grown epitaxially onto a silicon substrate, in the directionperpendicular to the substrate surface, by precision X-ray diffraction.As long as the critical stress for formation of misfit dislocations hasnot been exceeded, the difference in lattice parameter between theepitaxial layer and the substrate is proportional to the stress in thegrown layer. Another method is to measure the bow of a silicon substratebefore and after depositing the epitaxial Si--Ge--C layer. The change inbow is related by a simple formula to the stress in the Si--Ge--C layer.P. Singer, "Film Stress and How to Measure It", SemiconductorInternational, October 1992, p. 58, which is incorporated by referencein its entirety.

Advantages of Si--Ge--C having high etch selectivity:

As FIGS. 8 and 9 illustrate, the etch selectivity of the Si--Ge--Cepitaxial layer is much higher than that of Si--Ge--B in both KOH-H₂ Oand HNA etch solutions. In FIG. 8, the etch selectivity of Si--Ge--C in21 wt % KOH-H₂ O is plotted against carbon content in the epitaxiallayer, with the etch selectivity of a Si--Ge--B layer having 1E21 atomsper cm³ of Ge and 2E20 atoms per cm³ of B shown on the left side ofFIGS. 8 and 9 for comparison. The germanium content of thecarbon-containing etch-stop layers increases approximately in proportionto the carbon content to compensate strain, in atomic percent asfollows:

    ______________________________________    Carbon         Germanium Silicon    ______________________________________    2%             18%       80%    3%             27%       70%    4%             36%       60%    5%             42%       53%    ______________________________________

Here, etch selectivity in KOH-H₂ O is defined as the etch rate oflightly doped silicon divided by the etch rate of the etch-stop layer.FIG. 8 is a graph of the selectivity of Si--Ge--C and Si--Ge--B etchstops with respect to lightly doped silicon. A selectivity of 8000:1 in21 wt % KOH-H₂ O at 70° C. has been measured by employing black wax tomask the etch-stop layer before exposing to the etchant, measuring theresultant step in the etch-stop layer and comparing that distanceagainst the etch depth in lightly doped silicon under the same etchant.However, it is believed that by employing a silicon nitride mask, theselectivity of the Si₀.53 Ge₀.42 C₀.05 etch-stop will be even higher. Inany event, the selectivity of the Si--Ge--C etch-stop is apparentlyhigher than that of SiO₂ and is a significant improvement overSi--Ge--B, which shows a maximum selectivity of 2000:1 under the sameconditions.

FIG. 9 compares the etch selectivity of Si--Ge--B and Si--Ge--C as afunction of carbon concentration in five HNA solutions. For etch-stopremoval in HNA, etch selectivity is defined as the etch rate of theetch-stop layer divided by the etch rate of lightly doped silicon. FIG.9 illustrates that in HNA solution, the selectivity of the etch-stoplayer is enhanced by two innovations. The five HNA solutions areprepared from standard aqueous reagents which are as follows: 49 wt %HF, 70 wt % HNO₃, and 99 wt % CH₃ OOH. First, we discovered thatselectivity increases with carbon content, reaching a maximumselectivity at about 4 atomic percent carbon. Second, we discovered thatselectivity increases as the HNO₃ content of the HNA solution decreases,reaching maximum selectivity of more than 800:1 with the HF:HNO₃ :CH₃COOH volume ratio of 1:0.3:12 and 4 atomic percent carbon. A furtherbenefit of reducing the HNO₃ concentration in the HNA solution is thatthe solution is stable against the chemical reduction of HNO₃ to HNO₂, aproblem that plagues users of the conventional 1:3:8 and 1:3:12 HNAcompositions. It is an important advantage that our etch solutionsremain colorless and etch selectivity is independent of time for etchtimes of 30 minutes or more (1:3:8 and 1:3:12 HNA solutions turn yellowand lose selectivity rapidly over the first few minutes of use). Wediscovered the novel HNA formulations also will increase the etchselectivity of Si--Ge--B as the HNO₃ concentration of the solution isdecreased reaching a maximum of 50.

The unique properties of the Si--Ge--C etch-stop layers combined withthe new HNA etch compositions are suggested by the behavior of theSi--Ge--C surface in the etch solution. The etch solution is believed towork by simultaneously oxidizing the surface (performed by the HNO₃) andetching the oxide (performed by the HF). Si--Ge--C surfaces arehydrophilic when they emerge from the HNA solutions for all etchcompositions from 1:3:12 to 1:0.1:12, in contrast to lightly doped Siwith hydrophobic surfaces, indicating that the oxide removal performedby HF dominates the etching rate of Si--Ge--C layers.

Favorable results are believed to be obtained for etch selectivity inHNA when the silicon is in the range of 50 to 90 atomic percent, thegermanium is in the range of 9 to 50 atomic percent, and the carbon isin the range of 1 to 10 atomic percent, preferably in the range of 3 to6 atomic percent. More preferably, the Si--Ge--C is in the range for Si,Ge and C of 50-60 atomic percent, 36-45 atomic percent, and 4-5 atomicpercent, respectively.

We believe that the some of useful ranges of Si--Ge--C compositions are:silicon in a range of 0 to 99.9 atomic percent, preferred 45 to 83atomic percent, further preferred about 55.5 atomic percent, germaniumin a range of 1 to 92 atomic percent, preferred 15 to 50 atomic percent,and further preferred about 40 atomic percent, and carbon in the rangeof 0.1 to 10 atomic percent, preferred 2 to 6 atomic percent, andfurther preferred about 4.5 atomic percent.

Examples: Si--Ge--C Etch-Stop in Nanotechnology Applications

A. Membrane

One example of silicon membrane fabrication is illustrated in FIGS.10A-C. As shown in FIG. 10A, Si--Ge--C layer (101) is epitaxially grownonto a lightly doped, conventional polished silicon substrate (102).Favorable results can be achieved by using the processes describedearlier in the section on epitaxial growth of Si--Ge--C layers by CVD onsilicon. The silicon substrate (102) and Si--Ge--C layer (101) form astructure which is inverted as shown in FIG. 10B and an oxide layer(104) is formed on the back side of the silicon substrate (102). Apositive photoresist (not shown) is applied to the oxide layer (104),and then exposed through openings in the mask. The exposed and developedphotoresist is removed and then the underlying exposed oxide layer isremoved with a conventional HF solution. The exposed portion of the backside of the silicon substrate (102) is etched by a 21 wt % KOH-H₂ Oetchant that rapidly removes the lightly doped silicon substrate (102)as indicated by surface (103) formed during an intermediate stage of theetch, but then etches at a slower rate when it reaches the Si--Ge--C(101) as shown in FIG. 10C. As shown in FIG. 10C, when the etching iscomplete, the Si--Ge--C layer (101) remains as a membrane supported by asilicon substrate frame (105) defined by the remaining portions of thesilicon substrate (102). The carbon-to-germanium ratio in the etch-stoplayer is adjusted so that the Si--Ge--C membrane (101) is preferably intension. This ensures that the membrane (101) is "tight" and wellsupported by the frame (105).

B. Cantilever Beam

Another example of a micromechanical application of a Si--Ge--Cepitaxial etch-stop layer is in fabricating a cantilever beam asillustrated in FIGS. 11A-C. In FIG. 11A, a Si--Ge--C etch-stop layer(111) and a lightly doped silicon layer (112) are grown onto a siliconsubstrate (113). A masking oxide (114) shown in FIG. 11B is grown andpatterned using a technique like that used in the membrane fabrication,and the exposed portion of the lightly doped silicon (112 ) is etched bya 21 wt % KOH-H₂ O etchant that rapidly removes the lightly dopedsilicon but etches very slowly when it reaches the Si--Ge--C layer(111). Next, the Si--Ge--C layer (111) is etched by HNA that rapidlyremoves the exposed Si--Ge--C layer (111) as shown in FIG. 11C, but onlyetches lightly doped silicon very slowly. In this way, the cantileverbeam (115) formed from the lightly doped silicon layer (112) is freedfrom the surrounding silicon substrate. FIG. 11D is a perspectivedrawing showing the finished cantilever beam (115) attached to the restof the lightly doped silicon layer (112) at one end. Other precisionmicro-miniature structures can be fabricated using processes like this.

It will be appreciated by those of ordinary skill in the art that manyvariations in the foregoing embodiments are possible while remainingwithin the scope of the present invention. Thus, the invention hasapplications well beyond those enumerated. The present invention shouldthus not be considered limited to the preferred embodiments.

What is claimed is:
 1. A wet chemical etchant, comprising:3.5-15.8 molepercent of HF; 0.2-9.7 mole percent of HNO₃ ; 34.8-82.6 mole percent ofCH₃ COOH; 12.3-46.5 mole percent of H₂ O; andwherein the molecular ratioof HNO₃ to HF is less than about 0.78.
 2. A wet chemical etchant,comprising:6.2-8.2 mole percent of HF; 0.3-6.9 mole percent of HNO₃ ;53.8-71.1 mole percent of CH₃ COOH; 20.4-33.1 mole percent of H₂ O;andwherein the molecular ratio of HNO₃ to HF is less than about 0.78. 3.The wet chemical etchant of claim 1, wherein the molecular ratio of HNO₃to HF is less than 0.78.
 4. The wet chemical etchant of claim 1, whereinthe molecular ratio of HNO₃ to HF is less than about 0.20.
 5. The wetchemical etchant of claim 2, wherein the molecular ratio of HNO₃ to HFis less than 0.78.
 6. The wet chemical etchant of claim 2, wherein themolecular ratio of HNO₃ to HF is less than about 0.20.
 7. The wetchemical etchant of claim 1, wherein components are limited asfollows:7.3-9.5 mole percent of HF; 0.3-8.6 mole percent of HNO₃ ;62.9-79.0 mole percent of CH₃ COOH; and 11.2-21.3 mole percent of H₂ O.